High thermal diffusivity and high wear resistance tool steel

ABSTRACT

A tool steel family with outstanding thermal diffusivity, hardness and wear resistance has been developed, also exhibiting good hardenability. Also its mechanical strength, as well as its yield strength, at ambient and high temperature (superior to 600° C.) are high, due to a high alloying level in spite of the high thermal conductivity. Because of its high thermal conductivity and good toughness, steels of this invention have also good resistance to thermal fatigue and thermal shock. This steels are ideal for discontinuous processes where it is interesting to reduce cycle time and that require high hardness and/or wear resistance (plastic injection molding, other plastic forming processes and curing of thermosets, hot forming of sheet . . . ). These tool steels are also appropriate for processes requiring high wear resistance and good resistance to thermal fatigue (forging, hot stamping, light-alloy injection . . . ).

FIELD OF THE INVENTION

The present invention relates to a tool steel with very high thermaldiffusivity and high wear resistance, mainly abrasive. This tool steelalso shows good hardenability.

SUMMARY

Tool steels often require a combination of different properties whichare considered opposed. A typical example can be the yield strength andtoughness. For many metal shaping industrial applications in which thereis a heat extraction from the manufactured product which isdiscontinuous, thermal diffusivity is of extreme importance.Traditionally, for tool steels, this property has been consideredopposed to hardness and wear resistance. During plastic injection, hotstamping, even forging, metal injection, composite curing and othermetal shaping processes, wear resistance and high or very high thermaldiffusivity are often simultaneously required. For many of theseapplications, big cross-section tools are required, for whichhardenability of the material is also of extreme importance. Thermaldiffusivity (a) is related to other fundamental material properties likethe bulk density (ρ), specific heat (c_(p)) and thermal conductivity (λ)in the following way:

λ=ρ·c _(p) ·a

or if preferable:

a=λ/(ρ·c _(p))

Wear in material shaping processes is, primarily, abrasive and adhesive,although sometimes other wear mechanisms, like erosive and cavitative,are also present. To counteract abrasive wear hard particles aregenerally required in tool steels, these are normally ceramic particleslike carbides, nitrides, borides or some combination of them. In thisway, the volumetric fraction, hardness and morphology of the named hardparticles will determine the material wear resistance for a givenapplication. Also, the use hardness of the tool material is of greatimportance to determine the material durability under abrasive wearconditions. The hard particles morphology determines their adherence tothe matrix and the size of the abrasive exogenous particle that can becounteracted without detaching itself from the tool material matrix. Thebest way to counteract the adhesive wear is to use FGM materials(functionally graded materials), normally in the form of ceramic coatingon the tool material. In this case, it is very important to provide agood support for the coating which usually is quite brittle. To providethe coating with a good support, the tool material must be hard and havehard particles. In this way, for some industrial applications, it isdesirable to have a tool material with high thermal diffusivity at arelatively high level of hardness and with hard particles in the form ofsecondary carbides, nitrides and/or borides and often also primary hardparticles (in the case to have to counteract big abrasive particles).

Thermal gradients are the cause of thermal shock and thermal fatigue. Inmany applications steady transmission states are not achieve due to lowexposure times or limited amounts of energy from the source that causesa temperature gradient. The magnitude of thermal gradient for toolmaterials is also a function of their thermal conductivity (inverseproportionality applies to all cases with a sufficiently small Biotnumber).

Hence, in a specific application with a specific thermal flux densityfunction, a material with a superior thermal conductivity is subject toa lower surface loading, since the resultant thermal gradient is lower.The same applies when the thermal expansion coefficient is lower and theYoung's modulus is lower.

Traditionally, in many applications where thermal fatigue is the mainfailure mechanism, as in many casting or light alloy extrusions cases,it is desirable to maximize conductivity and toughness (usually fracturetoughness and CVN). Steels of the present invention prioritize wearresistance and diffusivity to CVN, although it is also considered veryimportant for some applications and, therefore, the intention is to tryto also maximize it but without renouncing to the other two properties.Usually, increasing the hardness of the tool steel will decrease bothtoughness and thermal diffusivity and will increase wear resistance. Agreater level of diffusivity for a given hardness level has beenachieved for steels of the present invention, usually together with agood hardenability and, for some cases, with an excellent toughnesscompromise.

For many applications thick tools are used, especially when sufficientstrength is required as for to require a thermal treatment. In thiscase, it is also very convenient to have a good hardenability to be ableto achieve the desired hardness level on surface and, preferably, allthe way to the nucleus. Hardenability is also very interesting for hotwork steels, since it is much easier to achieve high toughness with aquenched martensite structure than with a quenched bainite. Thus, thehigher the hardenability the less abruptly the quenching cooling willneed to be. A sudden cooling is more difficult to achieve and also moreexpensive and, since the forms of tools and components manufactured areoften complex, can lead to breaking of the parts being heat treated orsevere deformation.

Wear resistance and mechanical strength are often inversely proportionalto the toughness. Thus, it is not easy to get a simultaneous increase inboth properties. Thermal conductivity is a help in this case, since itallows a great increase of the thermal fatigue resistance, even if theCVN has been reduced to increase wear resistance or mechanical strength.

There are many other desirable properties, if not necessary, for hotwork steels that do not necessarily influence the longevity of the tool,but their production costs, like: ease of machining, welding or repairin general, support provided to the coating, costs . . . .

The authors have discovered that the problem to simultaneously obtainhigh thermal diffusivity, wear resistance and hardenability, togetherwith good levels of toughness, can be solved applying certain rules ofcomposition and thermo-mechanical treatments within the followingcompositional range:

% C_(eq) = 0.3-0.9 % C = 0.3-0.9 % N = 0-0.6 % B = 0-0.6 % Cr <2.8 % Ni= 0-3.8 % Si = 0-1.4 % Mn = 0-3 % Al = 0-2.5 % Mo = 0-10 % W = 0-12 % Ti= 0-2 % Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-4 % Nb = 0-1.5 % Cu = 0-2% Co = 0-6 % S = 0-1 % Se = 0-1 % Te = 0-1 % Bi = 0-1 % As = 0-1 % Sb =0-1 % Ca = 0-1,the rest consisting of iron and unavoidable impurities, wherein

% Ceq=% C+0.86*% N+1.2*% B

In the present invention it is always the case that:

% Mo+½·% W>3.0

Some of the selection rules of the alloy within the range andthermo-mechanical treatments required to obtain the desired high thermaldiffusivity to a high hardness level and wear resistance, are presentedin the detailed description of the invention section. Obviously, adetailed description of all possible combinations is out of reach. Thethermal diffusivity is regulated by the mobility of the heat energycarriers, which unfortunately can not be correlated to a singularcompositional range and a thermo-mechanical treatment.

In an additional aspect, the invention is related to a process tomanufacture a hot work tool steel, characterised in that the steel issubjected to a martensitic, bainitic or martensitic-bainitic quench withat least one tempering cycle at temperature above 590° C., so that asteel having a hardness above 47 HRc with a low scattering structurecharacterized by a diffusivity of 9 mm²/s or more is obtainable. Inanother embodiment, a steel having a hardness above 53 HRc with a lowscattering structure characterized by a diffusivity of 9 mm²/s or moreis obtainable. In an additional embodiment of this process, the steel issubjected to at least one tempering cycle at temperature above 660° C.,so that a steel having a hardness of 50 HRc or more with a lowscattering structure characterized by a diffusivity of 5.8 mm²/s or moreat 600° C. is obtainable.

STATE OF THE ART

Until the development of high thermal conductivity tool steels (EP1887096 A1), the only known way to increase thermal conductivity of atool steel was keeping its alloying content low and, consequently,showing poor mechanical properties, especially at high temperatures.Tool steels capable of surpassing 42 HRc after a tempering cycle at 600°C. or more, were considered to be limited to a thermal conductivity of30 W/mK and thermal diffusivity of 8 mm²/s and 6.5 mm²/s for hardnessabove 42 HRc and 52 HRc respectively. Tool steels of the presentinvention have a thermal diffusivity above 8 mm²/s and, often, above 12mm²/s for hardness over 52 HRc, and even more than 16 mm²/s for hardnessover 42 HRc, furthermore presenting a very good wear resistance and goodhardenability. Thermal diffusivity is considered the most relevantthermal property since it is easier to measure accurately and becausemost of the tools are used in cyclic processes, so that the thermaldiffusivity is much more important for evaluating performance of thetool than can be thermal conductivity.

Tool steels of the present invention have a wear resistance and hardnesshigher than steels described in EP2236639A1. The latter, on thecontrary, show a higher hardenability in the perlitic region and higherCVN compared to the tool steels with high thermal conductivity of thepresent invention. Hence, for applications where the main failuremechanism is thermal fatigue and no wear is present is better to usesteels of EP2236639A1 but, for applications where wear resistance isimportant, tool steels of this invention have great advantage.Furthermore, the steels of the present invention exhibit higher thermaldiffusivity for the same level of hardness. This is largely due to thefact that in EP2236639A1 carbides of the type of M₃Fe₃C, where Mcorresponds to Mo and/or W, are almost exclusively used, partly due tothe presence of % Ni in the matrix that penalizes the thermaldiffusivity in favour of hardenability, toughness (CVN) and lower linearthermal expansion coefficient. In the present invention there is lower %Ni and carbides are often partially replaced by harder carbides, evenwhen the elements forming harder carbides tend to be solubilized in theMo and/or W carbides, as is the case of % V.

The tool steels of the present invention can attain much higher levelsof thermal diffusivity than the tool steels of WO2004/046407 A1, wherethe high levels of % Cr impose very tight restrictions which are notobserved, on the compositions to be taken within the proposed range andthe small process window thereafter during the thermo-mechanicalprocessing to attain high levels of carrier mobility.

There are other inventions that may have compositional range overlap butdo not have anything to do with the present invention since rules forselecting the composition within the range and/or thermo-mechanicaltreatments required to achieve a structure with a matrix poor inelements in solid solution with great capacity to disperse heat energycarriers and having carbides with a high level of crystalline netperfection, and consequently a very low dispersion of heat energycarriers (mainly electrons and phonons), are not observed. This could bethe case of JP04147706 here the inventors, seeking an optimizedsuperficial oxide coating, are using levels of % Cr lower than thenormal ones (around 0.5%) to allow the mentioned oxidation with somespecific treatments at high temperature. In the present invention % Crhas the tendency to dissolve in the W and/or Mo carbides causing thedispersion of the heat energy carriers and thus their presence is alsoundesirable. This is the only point of coincidence that also, in thecase of JP04147706, does not lead to high thermal diffusivity in any ofthe examples described. At an even lower extent is the case ofJP11222650, where the inventors look for the presence of large amountsof primary carbides to resist massive wear as is the case for high speedsteel but with an exceptionally low content of % C to allow coldcoining.

Other cases may be misleading because of not making special mention orhaving a generic reference levels of non-functional elements for theapplication mentioned, this is often the case for % Cr, % Si and % Mn.In fact, it is difficult to achieve a low level for some elements insteels. For instance, a steel supposedly lacking Cr (0% Cr in nominalcomposition), especially if it is an alloyed steel, will probably have %Cr>0.3 if the steel is required, for some reason, to be made of selectedscrap. In the case where normal scrap can be used, significantlycheaper, a % Cr>0.5 would be expected. If, for a composition, the % Cris not mentioned then it means that its presence is not consideredimportant, but neither its absence. In this case, the content of % Crdoes not compel the use of especial scraps and, if there are not otherelements that require so, then a % Cr>0.5 can be expected. Even moreimportant is the placement of this % Cr, which will be predominantlydissolved in the carbides if no special measures are taken.

The case of % Si is slightly different, since it is possible to reduceits content through a refining process, such as ESR, although, due tothe narrow window of the process in this case, it is technologicallyvery difficult (and expensive, and therefore it is only carried out inthe case of seeking a specific functionality) to reduce the % Si below0.2 and, at the same time, to reach a low level of inclusions(especially oxides).

There are many tool steels having a composition with the potential ofachieving a high thermal conductivity and actually do not. This ismainly due to the two following reasons:

-   -   The thermo-mechanical treatments used do not pursue the        maximization of mobility of the heat energy carriers. Thermal        conductivity is not properly chosen as one of the main desirable        characteristic or, for materials previously developed, the        knowledge was lacking on how to attain a desired level of        thermal diffusivity before the publication of EP 1887096 A1, and        thus the phases present in the final microstructure are chosen        according to the optimization of some other properties desirable        for the application, generally a certain compromise of relevant,        to the application, mechanical properties. Thus, within a        composition, often strengthening mechanisms are chosen which are        very detrimental for thermal diffusivity.    -   In the melting, secondary metallurgy or re-melting process, not        enough attention is placed on what is happening beyond the        micrometric and nano-metric scales, and thus unfavorable atomic        scale arrangements take place, not necessarily in all phases        present, that lead to strong carrier scattering. Again this is        mainly due to the lack of knowledge before the publication of EP        1887096 A1.

There are several tool steels families that, with their nominal range ofcomposition, could have the potential to achieve high thermaldiffusivity when the correct strategy during the thermo-mechanicalprocess is employed according to the present application and EP 1887096A1, but do not end up with compositions capable of developing high orvery high thermal diffusivity. This is mainly due to the followingreasons:

-   -   The ratio of % C and that of carbide formers is not well        balanced to be able to minimize solid solutions in the metal        matrix, especially that of % C, and levels are provided that        cannot afterwards be properly managed by the thermo-mechanical        treatments used to pursue the maximization of mobility of the        heat energy carriers.    -   The nominal levels of certain critical elements are far away        from the real content values in the embodiment. For instance,        this is often the case for % Si and % Cr. While the nominal        composition can describe a certain level, especially in the case        of only upper bound descriptions, like % Cr<1 (or even without        mentioning the % Cr, which can lead to the erroneous assumption        that is 0%) and in the same fashion as often the case % Si<0.4,        it ends up by being % Cr>0.3 and % Si>0.25. This applies also        for all trace elements with a strong influence on the        conductivity of the matrix and even more those with a high        solubility in carbides and great potential for distortion of the        carbides structure. Usually, with the exception of % Ni and for        some applications the % Mn, no element is desirable in solid        solution with the matrix at a level higher than 0.5%.        Preferably, the percentage of these, individually in solid        solution, should not exceed 0.2%. If the main purpose of the        application is to maximize the thermal conductivity, then any        metallic element in solid solution with the matrix (obviously        including transition metals), with the exception of % Ni and in        some cases the % C and % Mn, should not exceed 0.1% or, even        better, 0.05%.

DETAILED DESCRIPTION OF THE INVENTION

To obtain tool steels with high thermal diffusivity and wear resistanceto high hardness levels with good hardenability, it has been observedthat, within the compositional range specified above, a number of rulesand general considerations in the selection of the composition withinthe range and the thermo-mechanical treatments to be used, some of whichare described below, have to be taken into account. Thermal diffusivityis a consequence of the scattering mechanisms on the phases present forall carrier types present. The perfection of the lattice plays animportant role, but also other scattering mechanisms are of relevance.In this document the thermal diffusivity itself will be used as ameasurement of the structure attained. Within a same chemicalcomposition different structures can be attained and thus also differentlevels of thermal diffusivity.

Tool steels of the present invention excel mainly because of their highthermal diffusivity and wear resistance. Wear resistance and toughnesstend to be inversely proportional, although different microstructuresreach different relationships, i.e., as a function of microstructuredifferent levels of toughness for the same elastic limit and hardness ata given temperature can be reached and, for a specific type of material,hardness tends to correlate with wear resistance unless the volumefracture or the morphology of wear resistant particles is significantlychanged. In this vein, it is well known that for most tool steels withmedium carbon content, pure microstructure of tempered martensite is theonly one that offers the best compromise of mechanical properties. Thismeans that it is important to avoid the formation of othermicrostructures like stable ferrite-perlite or metastable bainite duringcooling after the process of austenitization of the heat treatment.Therefore, fast cooling rates will be needed and, if higherhardenability is required, some alloying elements to delay the kineticsof the formation of these more stable structures should be used. Fromall possible alternatives those with less negative effects on thermaldiffusivity should be used.

A strategy to obtain wear resistance and higher elastic limit at hightemperatures and, at the same time, obtain high thermal conductivity, isthe use of carbides with high electron density, as secondary carbides ofthe M₃Fe₃C type and sometimes even primary carbides (M− should only beMo or W for a greater thermal conductivity). There are other carbidetypes (Mo, W, Fe) with high electron densities and with tendency tosolidify with a good crystalline perfection. Some elements like Zr andHf and, at a less extent, Ta for instance comparing to Cr, whendissolving with this type of carbides do not provide much distortion tothe crystalline structure and dispersion of charge carriers is small andso is the effect on thermal conductivity. Moreover, these high carbideforming elements tend to form separate MC type carbides, due to its highaffinity for C.

In fact, in the present invention it has been observed that the effectcan be quite positive if a moderate quantity of % V is used and it isbalanced with the presence of strong carbide former (preferably Zrand/or Hf). It has been seen that there can be amounts of % V up to 0.9with practically no formation of primary carbides (obviously dependingon the Ceq and the presence of other carbides, and for higher contentsof Ceq is necessary to reduce the percentage of V at a maximum of 0.8and even 0.5 or 0.4 to avoid the presence of primary carbides or massivedissolution in them) and with little dissolution in the carbides of (Fe,Mo, W), especially if used simultaneously with strong carbide formingelements, also there is a displacement of more carbon out of the matrixwith the consequent benefit to the overall thermal diffusivity (in thiscase, the benefit is remarkable with % Hf+% Zr+% Ta greater than 0.1,and very significant if it exceeds 0.4 or 0.6, depending on thequantities of % Ceq and % V present). In fact, this combination ishighly desirable as the percentage of V as the percentage of Zr, Hf andTa tend to significantly improve the wear resistance compared to a steelthat has only carbides (Fe, Mo, W), the same applied for % Nb. Theeffect becomes noticeable with % V=0.1 and remarkable with % V=0.3 or0.5, depending on the level of % Ceq. If extreme wear resistance withthe presence of primary carbides is to be achieved, as is the case inapplications with large abrasive particles such as in hot stamping ofuncoated sheet, then larger amounts of % V can be used, up to 1.5% oreven 2% is possible while maintaining a good level of thermaldiffusivity, especially if compensated with strong carbide formingelements. In this case, it can be convenient to have high levels ofstrong carbide forming elements combined with % V (% V+% Nb+% Hf+% Zr),above 1.2 or even 2.0 in weight percentage (for applications where agood wear resistance is needed, even 3.0, but then the cost of the alloyis increased). In this case, rarely any strong carbide forming element(% V, % Nb, % Ta, % Zr, % Hf) will individually exceed 3%, with theexception of % V where the upper limit is usually 4% in weight (forapplications where wear resistance is priorized at the expense of losingthermal diffusivity), or 1.8% for applications requiring very highthermal diffusivity and Nb that, due to its negative effect on thermaldiffusivity, tends to be used only to control grain size and when usedas primary carbide former will rarely be above 1.5%. It is desirable tohave most of the strong carbide formers bound in the carbides and notdissolved in the matrix, thus the level of % Ceq has to be finelyadjusted as explained later to minimize both the amount of strongcarbide formers and % Ceq in solid solution. As an example in mostapplications of this invention if % Ceq is smaller than 0.35 then % Vshould be kept below 1.7%. In general it is desired to mostly have Fe,Mo and W carbides (where obviously part of the C can be replaced by N orB), usually more that 60% and, optimally, more than 80% or even morethan 90% of these type of carbides. The dissolution of other metallicelements of these types of carbides (obviously in the case of carbidesmetallic elements are mainly transition elements) can exist, but it isdesired to be small to guarantee a high phononic conductivity. Normallyno other metallic element, apart from the principal Fe, Mo and W, shouldexceed 20% of the weight of all metallic elements of the carbide, forthis type of mainly desired carbides. Preferably should not be more than15% and even better a 5%. This is because they tend to form structureswith densities of solidification defects extremely low even for fastsolidification kinetics (therefore less structural elements to causedispersion of carriers).

As discussed before, the only exception is the presence of a limitedquantity of strong carbide forming elements, although the formation ofindependent carbides is preferable. In this case, Mo and W providesufficient obstacles for the formation of stable structures (perlite andferrite), although the formation of bainite is very fast. In some steelssuperbainitic structures can be formed applying a martempering heattreatment, which consists in the complete solubilisation of alloyingelements followed of a rapid cooling to a specific temperature (to avoidferrite formation) in the range of lower bainitic formation and anextended temperature maintenance to obtain a 100% bainitic structure.For the majority of the steels a pure martensitic structure isdesirable, so that in this system some bainitic transformation delayingelements must be added, since Mo and W are very inefficient in thisrespect. Generally, for this purpose % Cr is commonly used, but has anextremely negative effect on the thermal conductivity for this systembecause it dissolves in the M₃Fe₃C carbides and causes a greatdistortion, so it's much better to use strong carbide forming elementsand non soluble elements in carbides. These last elements will reducethe conductivity of the matrix and, thus, the ones with the minimumnegative effect should be used. Accordingly, the natural candidate isNi, but at the same time others can be used (a special mention should bemade to % B, for its effect with very low concentrations). Since in thepresent invention carbide formers with great affinity to C are tend tobe used for their positive effect on wear resistance, the necessary anddesirable quantity of delaying elements of the transformation kineticsto stable structures during quenching is lower. Usually, a quantity upto 1% in weight, and for large sections up to 3.0%, will be enough toget sufficient hardenability and contribute to the increase of toughnesswithout an excessive detriment of conductivity. Higher % Ni quantitiesprovide more toughness and a reduction in the linear thermal expandedcoefficient, but the priority of the present invention is a combinationof wear resistance with thermal diffusivity, thus, only for some specialapplications the strategy of using high contents of % Ni, with a maximumof 3.8%, can be used. There are applications where lower amounts of % Nialready lead to the desired effect, especially if the contents of % Mnand/or % Si are a bit higher (% Mn usually does not exceed 3%) orsections of the material used are smaller.

The use of % Mo as a single carbide former (obviously together with Fe),is advantageous when maximising thermal conductivity, but it has thedisadvantage to provide a higher thermal expansion coefficient and,thus, it decreases the thermal fatigue resistance. Hence it ispreferable to have a relation of 1.2 to 3 times more Mo than W, but notthe absence of W. The exception are the applications where only thermalconductivity is to be maximised together with toughness, but notparticularly thermal fatigue resistance. Hardenability and the alloycost, due to the high volatility of Mo and W prices, can lead tochanging preferences regarding the % W being the main element in %Mo_(eq), (where % Mo_(eq)=% Mo+½·% W).

High contents of Mo_(eq) can be used with high levels of C_(eq),resulting in an increased cost alloy, low toughness, very difficult toweld, complicated hardenability for large parts and limitedmachinability. But very high levels of wear resistance with good thermaldiffusivity can be achieved. For applications where the highlighteddrawbacks are not determinant these can be alloys of interest. This canbe the case for some cutting applications. Here, levels of C_(eq)usually superiors to 0.5% are used and, often, even over 0.6%. Levels of% Mo_(eq) are often above 5% and frequently above 6% or even 9%. Alsothe limits of the Mo_(eq)/C_(eq) ratio are shifted to superior levelscompared to the rest of the alloys of the present invention. Valueshigher than 16 are possible and, higher than 13, are probable.

In all document the term carbide is referring to the primary carbides aswell as the secondary, unless otherwise specified.

The more restrictive the % Si and % Cr the higher the thermalconductivity, although the solution is more expensive (also, someproperties, that could be important for some applications, and thuswould be desirable to be maintained, could get worse with the reductionof these elements below certain levels as, for instance, for toughnessdue to oxide inclusions in the case Al, Ti, Si and any otherdeoxidizing, are used in insufficient quantities or, in some cases ofcorrosion resistance, if % Cr or % Si are too low). Thus, often theremust be a compromise between increase of costs, toughness reduction,wear resistance or other relevant properties for certain applicationsand the benefit of higher thermal conductivity. Maximum thermalconductivity can be obtained only if levels of % Si and % Cr are below0.1% or, even better, if below 0.05%. To maximise thermal diffusivity,also levels of the rest of the elements, with the exception of % C, %Mo, % W, % V, % Zr, % Hf, % Ta, % Nb and in some instances % Mn and %Ni, must be as low as possible (below 0.05 is technically possible withan acceptable cost for most of the applications, although a maximum of0.1 is, of course, less expensive). For some applications in whichtoughness is especially important less restrictive levels of % Si mustbe employed (is the least detrimental to the thermal conductivity of alliron deoxidants elements) and thus give up to some thermal conductivity,to ensure the inclusion level is not too high. Depending on the levelsof % C, % Mo and % W used, there can be sufficient hardenability,especially in the perlitic zone. For cases of large components, where itis not possible to avoid the formation of bainite during quenching, theuse of elements in solid solution to prevent the formation of coarsecementite precipitates (Fe₃C) that entail very low toughness, such as %Al and % Si, may be interesting. Generally below 0.4, exceptionally withlevels of around 1% and, very exceptionally, above 2% and for the % Al,up to a maximum of 2.5%. The levels of % Mo, % W and % C used to obtainthe desired mechanical properties must be balanced to achieve a highthermal conductivity, so that within the matrix remain the least amountof these elements in solid solution. The same applies for the rest ofcarbide formers that could be used to obtain a certain tribologicalresponse (like % V, % Zr, % Hf, % Ta, . . . ).

For some applications some environmental resistance can be of interestand, thus, be desirable to have some % Cr or % Si in solid solution(oxidation resistance to high temperature). The negative effect onthermal diffusivity can be moderated through carbon fixing with strongercarbide formers elements. Without the later, % Cr should not exceed 2%and, preferable, 1.5%. Although in the presence of V, % Nb, % Ta, % Zrand % Hf, and preferably the last two or three, levels close to 3% of Crcan be achieved maintaining a good thermal diffusivity, and even 1.4%for the case of Si. In fact for most applications % Cr<2.8% is requiredif the thermal diffusivity needs to be high. Many compositions require %Cr<2.5% to be able to attain high thermal diffusivity with the properthermo-mechanical processing (which is composition dependent, asexplained). At this level the environmental protection effect is onlysomewhat noticeable if the % Cr is mainly left in solid solution in thematrix. Finally, a much greater range of compositions can attain highthermal diffusivity when the proper thermo-mechanical treatment isapplied, if % Cr is restricted to remain below 1.9%.

The simplest compositional rule to describe the compositions within therange that are capable of attaining a high thermal diffusivitysimultaneously to a high wear resistance can be based on a ratioR=Mo_(eq)/C_(eq), where % Mo_(eq)=% Mo+½·% W and % C_(eq)=% C+0.86*%N+1.2*% B. This rule applies only for big enough contents of % C_(eq)(normally 0.32 min, preferably 0.35 min and most accurately when 0.38minimum % Ceq) and % Mo_(eq) (normally 3.2 min, preferably 3.4 min andmost accurately when 3.6 minimum % Moeq). It is also a rule that canonly be used for lower % Cr contents, normally % Cr<2.5%, and desirably% Cr<1.9%. The minimum value for R results when computing the % Mo_(eq)minimum for the rule to apply divided by 0.9 which is the maximum %C_(eq) for the present invention (for example for a minimum Moeq=3.2then the minimum R value results to be 3.56). The maximum value for Rhas been observed to be possibly 11.5, preferably 10.8 and optimally10.5 for low % C_(eq) values. Low % Ceq values are for this rule thoseunder 0.35%, occasionally under 0.36% or even under 0.37%. For high %C_(eq) values the maximum value for R has been observed to be possibly16.8, preferably 16.0 and optimally 15. High % C_(eq) values are forthis rule those above 0.38%, occasionally above 0.40% or even above0.45%. For intermediate values of % C_(eq), the maximum value for R hasbeen observed to be 14, preferably 13, and optimally 12.

Generally, to solely maximize thermal diffusivity (i.e. there are notother properties of great importance), it is convenient to observe thefollowing alloying rule (to minimize the % C in solid solution), if atempered martensite or bainite microstructure withstanding mechanicalrequirements wants to be obtained. The formula must be corrected ifcarbide formers with high affinity for the % C (like Hf, Zr or Ta, evenNb) are used. It must be also modified if % Cr>0.2 or Mo_(eq)>7:

0.02<xC_(eq)−solC−AC·[(xMo−solMo)/(3·AMo)+(xW−solW)/(3·AW)+(xV−solV)/AV]>0.265

where:xC_(eq)—carbon weight percentage;xMo—molybdenum weight percentage;xW—tungsten weight percentage;xV—vanadium weight percentage;AC—carbon atomic mass (12.0107 u);AMo—molybdenum atomic mass (95.94 u);AW—tungsten atomic mass (183.84 u);AV—vanadium atomic mass (50.9415 u);solC—carbon percentage in solid solution;solMo—molybdenum percentage in solid solution;solW—tungsten percentage in solid solution;solV—vanadium percentage in solid solution.

For an even higher thermal conductivity it is even more desirable tohave:

0.04<xC_(eq)−solC−AC·[(xMo−solMo)/(3·AMo)+(xW−solW)/(3·AW)+(xV−solV)/AV]>0.22

And still better:

0.09<xC_(eq)−solC−AC·[(xMo−solMo)/(3·AMo)+(xW−solW)/(3·AW)+(xV−solV)/AV]>0.18

To compensate for the presence of other % C avid carbide formers, anextra term must be added to the formula for each type of % C avidcarbide former:

−AC*xM/(R*AM)

where:xM—carbide former weight percentage;AC—carbon atomic mass (12.0107 u);R—number of carbide former units per carbide unit (for example: 1 if thecarbide type is MC, 23/7 if the carbide type would be M₂₃C₇ . . . )AM—carbide former atomic mass.

This balance provides an extraordinary thermal conductivity if thereinforcing ceramic particles formers, including the non metallic part(% C, % B and % N), are taken into the carbides (as an alternativenitrides, borides and intermediate substances). Then, the appropriatedthermal treatment must be applied. This thermal treatment will have aphase in which most of the elements will be dissolved (austenitizationto sufficiently high temperature, usually around 1080° C. for moderatedMo_(eq) levels, 1120° C. for medium levels of Mo_(eq) and 1240° C. forhigh levels of Mo_(eq), exceptionally, if distortion of the heattreatment is of great importance for the application, loweraustenitization temperatures can be used). An abrupt cooling willfollow, its intensity will be determined by the desired mechanicalproperties, although stable structures should be avoided since phaseswith big quantities of % C and carbide formers in solid solution areimplied. Metastable microstructures are even worse, since themicrostructure distortion caused by carbon is even greater, hencethermal conductivity is lower, although once these metastable structureshave relaxed the carbide formers place themselves in the desiredposition. Martensite and bainite tempered following this procedure willbe the desired microstructures for this case. The largest possiblecarbide substitution of Fe by Mo, W and all carbide forming elementswith greater affinity for carbon other than Cr are desired, so thetempering strategy selected has a great influence in the final thermalconductivity, with particular relevance to the final temperingtemperature and minimum tempering temperature. For hardness over 40 HRc,the highest possible temperature is desirable for the last tempering ifthermal diffusivity is to be maximized, and this approach is used to setthe intermediate tempering strategy. That is, the same final hardnesslevel can be achieved with different sequences of tempering and the oneusing a higher final tempering temperature is chosen, if the onlyobjective is to maximize the thermal diffusivity at a certain level ofhardness. So, usually, unusually high final tempering temperatures endup being used, often above 600° C., even when hardness over 50 HRc arechosen. In steels of the present invention it is usual to achievehardness of 47 HRc, even more than 52 HRc, and often more than 53 HRcand with the embodiments regarded as particularly advantageous due totheir wear resistance, hardness above 54 HRc, and often more than 56 HRcare possible with even one tempering cycle above 590° C., giving a lowscattering structure characterized by a thermal diffusivity greater than8 mm²/s and, generally, more than 9 mm²/s, or even more than 10 mm²/s,when particularly well executed then greater than 11 mm²/s, even greaterthan 12 mm²/s an occasionally above 12.5 mm²/s. As well as achievinghardness greater than 42 HRc, even more than 50 HRc with the lasttempering cycle above 600° C., often above 640° C., and sometimes evenabove 660° C., presenting a low scattering structure characterized by athermal diffusivity higher than 10 mm²/s, or even than 12 mm²/s, whenparticularly well executed then greater than 14 mm²/s, even greater than15 mm²/s and occasionally above 16 mm²/s. Those alloys can present evenhigher hardness with lowering tempering temperatures, but for most ofthe intended applications a high tempering resistance is very desirable.As can be seen in the examples with some very particular embodimentswith high carbon and high alloying, leading to a high volume fraction ofhard particles, hardness above 60 HRc with low scattering structurescharacterized by thermal diffusivity above 8 mm²/s and generally morethan 9 mm²/s are possible in the present invention.

To attain the high levels of hardness and wear resistance oftendesirable in the present invention, considerably high levels of thevolume fraction of hard particles have to be used. The volume fractionof hard particles (carbides, nitrides, borides and mixtures thereof) isoften above a 4% preferably above a 5.5% and for some high wearapplications, even above a 9%. Size of primary hard particles is veryimportant to have an effective wear resistance and yet not excessivelysmall toughness. The inventors have observed that for a given volumefraction of hard particles overall resilience of the material diminishesas the size of the hard particles increases, as would be expected. A bitmore surprisingly it has also been observed that when the size of hardparticles is increased, the overall fracture toughness increases if thefracture toughness of the particles themselves is maintained. When itcomes to abrasive wear resistance it has been observed the existence ofa critical hard particle size, below which the hard particle is noteffective against the abrasive agent. This critical size depends on thesize of the abrasive agent and the normal pressure. For someapplications where the abrasive particles are of small size (normallybelow 20 microns), it can be desirable to have primary hard particlessmaller than 10 microns or even smaller than 6 microns, but in any casewith an average size not smaller than 1 micron. For applications wherebig abrasive particles cause the wear, big primary hard particles willbe desirable. Therefore, for some applications it is desirable to havesome primary hard particles bigger than 12 microns, often greater than20 microns and for some particular applications even greater than 42microns.

For applications where mechanical strength more than wear resistance areimportant, and it is desirable to attain such mechanical strengthwithout compromising all too much toughness, the volume fraction ofsmall secondary hard particles is of great importance. Small secondaryhard particles, in this document, are those with a maximum equivalentdiameter (diameter of a circle with equivalent surface as the crosssection with maximum surface on the hard particle) below 7.5 nm. It isthen desirable to have a volume fraction of small secondary hardparticles for such applications above 0.5%. It is believed that asaturation of mechanical properties for hot work applications occurs ataround 0.6%, but it has been observed by the inventors that for someapplications requiring high plastic deformation resistance at somewhatlower temperatures it is advantageous to have higher amounts than these0.6%, often more than 0.8% and even more than 0.94%. Since themorphology (including size) and volume fraction of secondary carbideschange with heat treatment, the values presented here describeattainable values with proper heat treatment.

Cobalt has often been used in hot work tool steels principally due tothe increase in mechanical strength, and in particular the increase ofyield strength maintained up to quite high temperatures. This increasein yield strength is attained trough solid solution and thus it has aquite negative effect in the toughness. The common amounts of Co usedfor this propose is 3%. Besides the negative effect in toughness it isalso well known the negative effect in the thermal conductivity. Theinventors have seen that within the compositional ranges of the presentinvention it is possible to use Co, and attain an improved yieldstrength/toughness relation since Co can promote the nucleation ofsecondary hard particles and thus keep their size small. It has alsobeen seen that for some compositions of the present invention, whenadding Co the Thermal diffusivity does indeed decrease at roomtemperature, but then can actually increase at higher temperatures(normally above 400° C.) if the correct thermo-mechanical treatment isapplied. The inventors have seen that the best results are encounteredwhen % Co is above 1.3%, preferably above 1.5% and optimally above 2.4%.Also % C should exceed 3.2%, preferably 3.4% and optimally 3.6%. Ifthermal conductivity at high temperatures is of outmost importance forthe application a special care has to be taken not to have excessive %V, it should be kept below 2.8%, preferably below 2.3% and optimallybelow 1.7%. Finally % Moeq should normally exceed 3.3% often 3.5% andeven 4.0%. Heat treatment has to be selected with a rather highaustenitization temperature and an abnormally high temperingtemperatures, actually more than 55 HRc commonly achieved with at leastone tempering cycle at 630° C. or even above, 50 HRc can be maintainedeven with one tempering cycle at 660° C. or more. Properthermo-mechanical processing together with the compositional rules justexplained have to be implemented to minimize scattering at hightemperatures, the optimized arrangements is characterized by providingdiffusivities of more than 5.8 mm2/s, often more than 6.1 mm2/s and evenmore than 6.5 mm2/s at measuring temperatures as high as 600° C.

When mainly remaining in the carbide system Mo_(x)W_(3-x)Fe₃C, one ofthe preferred ways to balance the contents of % W, % Mo and % C in thepresent invention is through the adhesion to the following alloyingrule:

% C_(eq)=0.4+(% Mo_(eq(real))−4)*0.04173

where:

Mo_(eq(real))=% Mo+(AMo/AW)*% W.

with:AMo—molybdenum atomic mass (95.94 u);AW—tungsten atomic mass (183.84 u);so that, at the end:

Mo_(eq(real))=% Mo+0.52*% W.

If the expression is normalized in a parameter K=% C_(eq)/(0.4+(%Mo_(eq(real))−4)*0.04173), the desirable values for this parameter, forthe present invention, are as follows:

It has been observed that when carbon content is low (that is to say %Ceq<0.39, preferably % Ceq<0.36 and optimally % Ceq<0.35), the parameterK should exceed 0.75, preferably 0.76, more preferably 0.86 andoptimally 0.88. In fact for some embodiments for applications requiringvery high wear resistance, K will normally be higher than 0.92. A verygood performance will be obtained as already described, at the expenseof a higher cost, when adding elements that strongly bond carbon to thecarbides. In the case here dealing with low % Ceq it is especiallydesirable that the added amount of % Hf, % Zr, % Ta and % Nb exceed0.07%, preferably 0.09% and optimally 0.1%. Given that Nb can be quitedetrimental for the thermal diffusivity for some applications it willnot be desirable (% Nb<0.09) and then the contents of Hf, Zr and Ta inthe sum should exceed 0.01%, preferably 0.07% and in applicationsrequiring high wear resistance with very high thermal diffusivity andwhere Zr is chosen as the main former of hard carbides, then contentsabove 0.14%, preferably above 0.2% and even above 0.4%, will bedesirable. In these cases with the presence of strong carbon binders therestrictions on K can be relaxed around a 3% to 5% for alloys with lowcarbon content as hereby described.

On the other hand, when carbon content is not low (that is to say %Ceq>=0.39, preferably % Ceq>=0.36 and optimally % Ceq>=0.35), theparameter K should exceed 0.6, preferably 0.75, more preferably 0.84 andoptimally 0.87. In this case if elements are used that strongly bondcarbon (nitrogen or boron) to the carbides (nitrides, borides ormixtures) in the fashion described in the last paragraph, then therestrictions on K can be relaxed very severely, for some applicationseven eliminated.

The authors have observed that good combinations of wear resistance andthermal diffusivity can be obtained for very high values of K if allother alloying and thermal processing rules are observed, normally intheir most stringed version, but of course the best results are obtainedwhen K does not exceed 3, preferably 1.5 and optimally 1.3.

An especially interesting embodiment, when the main goal for the chosenapplication is the maximization of the thermal diffusivity to thehighest possible level of hardness, arises when applying this alloyingrule together with very low levels of % Cr, especially in dissolutionwith the carbides, as described above.

It has also been observed by the authors that it is possible to attainconsiderably high thermal diffusivity and wear resistance when usingmuch higher levels of % Mo and % W than described in the last coupleparagraphs. The level of thermal diffusivity for a given hardness levelcannot be optimized to such high values as when applying the previouslydescribed alloying rules. On top this comes at a considerably highercost, so obviously is not the preferred way for most applications, butin can be advantageous for some very concrete cases. For example if aspecial oxidation color is desirable, or when ferrite/perlitehardenability wants to be extended and the usage of other much moreeffective elements is nor recommendable. In such case the parameter Khas to be selected to be quite low, indeed it should be lower than 0.81,preferably lower than 0.79 and optimally lower than 0.75. This has tohappen for large enough values of % Ceq, normally larger than 0.33, evenlarger than 0.35 and occasionally larger than 0.41.

The teachings of this inventions can be applied to the describedcompositional range for alloys with a % Moeq>3.0%. To be more precise itcan be described in terms of % Moeq(real), in which case for mostapplications the teachings work for values superior to 3.3%, and evenmore generalized in terms of applications for values of %Moeq(real)>3.6% and when % Moeq(real)>3.8% then the density ofcompositions which can attain a high thermal diffusivity and wearresistance within the range is significantly greater, and covers mostapplications (one exemption is for example applications withexceptionally high hardness or wear resistance). In the same way when itcomes to % Ceq, while the teachings of the present invention workalready for values higher than 0.31%, when % Ceq>0.33%, and even morefor % Ceq>0.36% the density of compositions which can attain a highthermal diffusivity and wear resistance within the range issignificantly greater, and covers most applications (one exemption isfor example applications with exceptionally high hardness or wearresistance).

To increase machinability S, As, Te, Bi or even Pb, Ca, Cu, Se, Sb orothers can be used, with a maximum content of 1%, with the exception ofCu, than can even be of 2%. The most common substance, sulfur, has, incomparison, a light negative effect on the matrix thermal conductivityin the normally used levels to increase machinability. However, itspresence must be balanced with Mn, in an attempt to have everything inthe form of spherical manganese bisulphide, less detrimental fortoughness, as well as the least possible amount of the remaining twoelements in solid solution in case that thermal conductivity needs to bemaximized.

Another hardening mechanism can be used in order to search for somespecific combination of mechanical properties or environmentaldegradation resistance. It is always the intention to maximize thedesired property, but trying to have minimal possible adverse impact onthermal conductivity. Solid solution with Cu, Mn, Ni, Co, Si, etc. . . .(including some carbide formers with less affinity to carbon, like Cr)and interstitial solid solution (mainly with C, N and B). For thispurpose, precipitation can also be used, with an intermetallic formationlike Ni₃Mo, NiAl, Ni₃Ti . . . (also of Ni and Mo, small quantities of Aland Ti can be added, but special care must be taken for Ti, since itdissolves in M₃Fe₃C carbides and a 2% should be used as a maximum).Finally, other carbide types can also be used, but it is usuallydifficult to maintain high levels of thermal conductivity, unlesscarbide formers present a very high affinity with carbon, as it has beendescribed throughout this document. Co can be used as a hardener bysolid solution or as a catalyst of Ni intermetallic precipitation,rarely in contents higher than 6%. Some of these elements are also notas harmful when dissolved in M₃Fe₃C carbides, or other carbides of (Fe,Mo, W), this is specially the case for Zr and Hf and, to a lesserextent, for Ta, these can also limit V and Nb solubility.

When amounts are measured in weight percentage, atomic mass and theformed type of carbide determine if the quantity of a used elementshould be big or small. So, for instance, 2% V is much more than 4% W. Vtends to form MC carbides, unless it dissolves in other existingcarbides. Thus, to form a carbide unit only a unit of V is needed, andthe atomic mass is 50.9415. W tends to form M₃Fe₃C carbides in hot worksteels. So three units of W are needed to form a carbide unit, and theatomic mass is 183.85. Therefore, 5.4 more times carbide units can beformed with 2% V than with 4% W.

Tool steel of the present invention can be manufactured with anymetallurgical process, among which the most common are sand casting,lost wax casting, continuous casting, melting in electric furnace,vacuum induction melting. Powder metallurgy processes can also be usedalong with any type of atomization and eventually subsequent compactingas the HIP, CIP, cold or hot pressing, sintering (with or without aliquid phase and regardless of the way the sintering process takesplace, whether simultaneously in the whole material, layer by layer orlocalized), laser cusing, spray forming, thermal spray or heat coating,cold spray to name a few of them. The alloy can be directly obtainedwith the desired shape or can be improved by other metallurgicalprocesses. Any refining metallurgical process can be applied, like VD,ESR, AOD, VAR . . . . Forging or rolling are frequently used to increasetoughness, even three-dimensional forging of blocks. Tool steel of thepresent invention can be obtained in the form of bar, wire or powder(amongst others to be used as solder or welding alloy). Even, a low-costalloy steel matrix can be manufactured and applying steel of the presentinvention in critical parts of the matrix by welding rod or wire madefrom steel of the present invention. Also laser, plasma or electron beamwelding can be conducted using powder or wire made of steel of thepresent invention. The steel of the present invention could also be usedwith a thermal spraying technique to apply in parts of the surface ofanother material. Obviously the steel of the present invention can beused as part of a composite material, for example when embedded as aseparate phase, or obtained as one of the phases in a multiphasematerial. Also when used as a matrix in which other phases or particlesare embedded whatever the method of conducting the mixture (forinstance, mechanical mixing, attrition, projection with two or morehoppers of different materials . . . ).

Tool steel of the present invention can also be used for themanufacturing of parts under high thermo-mechanical loads and wearresistance or, basically, of any part susceptible to failure due to wearand thermal fatigue, or with requirements for high wear resistance andwhich takes advantage of its high thermal conductivity. The advantage isa faster heat transport or a reduced working temperature. As an example:components for combustion engines (such as rings of the engine block),reactors (also in the chemical industry), heat exchange devices,generators or, in general, any power processing machine. Dies forforging (open or closed die), extrusion, rolling, casting and metalthixoforming. Dies for plastic forming of thermoplastics and thermosetsin all of its forms. In general, any matrix, tool or part can benefitfrom increased wear resistance and thermal fatigue. Also dies, tools orparts that benefit from better thermal management, as is the case ofmaterial forming or cutting dies with release of large amounts of energy(such as stainless steel or TRIP steels) or working at high temperatures(hot cutting, hot forming of sheet).

Additional embodiments are described in the dependent claims.

EXAMPLES

Some examples indicate the way in which the steel composition of theinvention can be specified with higher precision for different hotworking applications:

Example 1

Dies for the stamping or press hardening of sheet. In this case maximumpossible thermal diffusivity is desired at high hardness. The desiredwear resistance depends on the sheet coating.

-   -   Sheets coated with Zn, AlSi or other inorganic coatings (the        same compositions are optimized for the manufacture of injection        molds for thermoplastics, especially when steels described below        are made by powder metallurgy):

For this purpose in the context of the present invention the followingcompositional range can be used:

C_(eq): 0.3-0.6 Cr<3.0% (preferably Cr<0.1%)V: 0-0.9% (preferably 0.3-0.8%)Si: <0.15% (preferably % Si<0.1, but with an acceptable level of oxideinclusions)

Mn: <0.5% Mo_(eq): 3.5-5.5

where

Mo_(eq)=% Mo+½% W and

C_(eq)=% C+0.86*% N+1.2*% B

The rest of the elements should be kept as low as possible and, in anycase, always be below 0.15%, with the exception of strong carbideformers (% Ta, % Zr, % Hf). All values are given in weight percentage.

The following three examples show properties that can be obtained:

Hardness Therm. Diff. % C % Mo % W % V % Cr % Si % Mn Other HRc mm²/s0.40 3.6 1.4 0.3 <0.01 <0.05 <0.01 — 52 11.47 0.45 1.6 4.5 0.4 <0.01<0.05 <0.01 — 52-53 10.96 0.41 3.5 1.4 0.8 1.3 <0.05 <0.01 — 50 9.32 *In all cases heat treatment which maximizes diffusivity at the indicatedhardness has been applied, minimizing the presence of elements insolution with the matrix, except for % Cr, especially minimizing thepresence of % C and, to a lesser extent, % V in the matrix. In all casesthis means very high austenitization temperatures, from 3 to 5 temperingcycles, with the latest in the range 600-640° C.

An advanced optimization is obtained when elements strongly reactingwith % C to form carbides (also % N and % B) are employed. Severalexamples show the properties that can be obtained:

Max. Hard./ Therm. Tem. T° Hardness diff. % C % Mo % W % V % Cr % Si %Mn Other HRc/° C. HRc mm²/s 0.50 3.6 1.4 0.5 <0.01 <0.05 <0.01 Hf, Zr,Nb 56-57 12.73 0.50 3.6 1.4 0.5 <0.01 <0.05 <0.01 Hf, Zr, Nb 54 13.930.33 3.36 1.91 <0.01 <0.01 <0.05 0.4 Hf, Zr, Nb, 50 13.04 B = 0.16 0.333.36 1.91 <0.01 <0.01 <0.05 0.4 Hf, Zr, Nb, 43 16.62 B = 0.16 0.42 3.451.4 0.6 2.2 <0.05 <0.01 Hf = 0.3, 52 10.42 Zr = 0.2 0.36 3.67 1.33 0.46<0.01 <0.05 <0.01 Hf, Nb, 56/600 54 12.83 Zr = 0.25 0.36 3.75 1.34 0.5<0.01 <0.05 <0.01 Zr, Nb, 57/600 54.5 13.01 Hf = 0.28 0.32 3.67 1.670.23 <0.01 <0.05 <0.01 Zr = 0.22 54/615 53.5 12.13 Hf = 0.42 0.33 3.81.22 0.40 <0.01 <0.05 <0.01 Hf, Zr, Nb 55/610 42 16.01 0.38 3.74 1.360.02 <0.01 <0.05 <0.01 Hf, Nb, 54/605 51.5 13.34 Zr = 0.55 0.38 3.741.36 0.02 <0.01 <0.05 <0.01 Hf, Nb, 54/605 44 16.04 Zr = 0.55 0.36 3.661.26 0.01 <0.01 <0.05 <0.01 Hf, Nb, 53/605 51.5 12.10 Zr = 0.44 * In allcases heat treatment which maximizes diffusivity at the indicatedhardness has been applied, minimizing the presence of elements insolution with the matrix, except for % Cr, especially minimizing thepresence of % C and, to a lesser extent, % V in the matrix. In all casesthis means very high austenitization temperatures, from 3 to 5 temperingcycles, with the latest in the range 610-680° C. Being % Hf: 0.10-0.22,% Zr: 0.05-0.18 y % Nb: about 0.07, unless specifically indicated.

-   -   Uncoated sheets and, therefore, with iron oxides that can be        large:

For this purpose, in the context of the present invention, the followingcompositional range can be used:

C_(eq): 0.4-0.9 Cr<3.0% (preferably Cr<0.1%)V: 0-2.0% (preferably 0.4-0.8%)

Si: <0.5% Mn: <1.0% Mo_(eq): 3.5-9

where

Mo_(eq)=% Mo+½% Wy

C_(eq)=% C+0.86*% N+1.2*% B

The rest of the elements should be kept as low as possible and, in anycase, always be below 0.15%, with the exception of strong carbideformers (% Ta, % Zr, % Hf). All values are indicated in weightpercentages.

The following examples show the properties that can be obtained:

Hardness Therm. diff. % C % Mo % W % V % Cr % Si % Mn Other HRc mm²/s0.60 3.6 1.2 0.62 <0.01 0.14 0.54 — 58 11.17 0.60 3.6 1.2 0.62 <0.010.14 0.54 — 47.5 12.47 0.85 6.48 4.0 <0.01 0.02 0.2 0.22 — 52.5 10.960.85 6.48 4.0 <0.01 0.02 0.2 0.22 — 45 12.87 0.79 6.42 3.78 0.41 <0.010.1 0.1 Hf, Zr 54 11.42 0.79 6.42 3.78 0.41 <0.01 0.1 0.1 Hf, Zr 4612.38 0.45 3.87 1.67 0.49 <0.01 0.45 <0.01 — 51 10.65 *In all cases heattreatment which maximizes diffusivity at the indicated hardness has beenapplied, minimizing the presence of elements in solution with thematrix, especially minimizing the presence of % C and, to a lesserextent, % V in the matrix. In this case, also seeking the highestpossible presence of primary carbides. In all cases this means very highaustenitization temperatures, from 2 to 4 tempering cycles, with thelatest in the range 550-620° C. Being % Hf: 0.10-0.22, % Zr: 0.05-0.18 y% Nb: about 0.07, unless specifically indicated.

Example 2

For closed-die forging. In this case, a simultaneous optimization ofwear resistance and thermal fatigue has to be achieved, therefore,maximum thermal diffusivity and wear resistance are desirable (presenceof primary carbides) maintaining also maximized CVN. For dies or largeparts subject to thermal shock or thermal fatigue a good CVN should bemaintained, even when the treatment cannot be fully martensitic, inwhich case Si or Al are used to hinder the precipitation of thickcementite (Fe₃C), or % Ni is used to improve the hardenability in theferritic-perlitic zone and decrease the linear thermal expansioncoefficient. In this case, tool steels in the following range can beused (powder metallurgy steels except for applications where the presentabrasive particles are very large). Steels of the present invention areparticularly attractive for applications where wear is the predominantfailure mechanism:

For this purpose, in the context of this invention, a compositionalrange of the following type can be used:

C_(eq): 0.3-0.6 Cr<0.1% (preferably Cr<0.05%)

Si: <1.4% Al: 0-2% Mn: <1.5% Mo_(eq): 3.0-7.0

where

Mo_(eq)=% Mo+½% W

The rest of the elements should be kept as low as possible and, in anycase, always be below 0.15%, with the exception of strong carbideformers (% Ta, % Zr, % Hf). All values are given in weight percentages.

Five examples show the properties that can be obtained:

Therm. diff. Hardness mm²/s at % C % Mo % W % V % Cr % Si % Mn Other HRc400° C. 0.37 3.3 1.01 <0.01 <0.01 <0.05 <0.01 Hf, Zr, 45.5 11.14 Ni =2.9 0.31 3.08 0.86 <0.01 <0.01 <0.05 0.16 Hf, Zr, 44 12.69 Ni = 2.3 0.53.65 1.27 0.45 <0.01 0.1 <0.01 Al = 0.7 53 10.12 0.5 3.73 1.52 0.17<0.01 0.8 <0.01 Hf, Zr, S 51 9.74 0.53 3.61 1.35 0.44 <0.01 <0.05 0.6 Al= 0.8 55 9.62 * In all cases heat treatment which maximizes diffusivityat the indicated hardness has been applied, minimizing the presence ofelements in solution with the matrix, especially minimizing the presenceof % C and, to a lesser extent, % V in the matrix. In this case, alsoseeking the highest possible presence of primary carbides. In all casesthis means very high austenitization temperatures, from 3 to 5 temperingcycles, with the latest in the range 590-660° C. Being % Hf: 0.10-0.22,% Zr: 0.05-0.18 y % Nb: about 0.07, unless specifically indicated.

Example 3

Some closed-die forging applications, require predominantly yieldstrength at high temperatures, good toughness, specially fracturetoughness and CVN, and as good as possible wear resistance. When thecontact times are long, or the temperature of the forged piece high,thermal diffusivity at high temperatures and good tempering resistanceare of outmost importance. In this case the correct usage of % Co isvery important. For this purpose, in the context of this invention, acompositional range of the following type can be used:

C_(eq): 0.32-0.7 V: <2.8% Si: <1.4% Mn: <1.5% Co: 1.3-6% Mo_(eq):3.3-7.0

where

Mo_(eq)=% Mo+½% W

The rest of the elements should be kept as low as possible and, in anycase, always be below 0.15%, with the exception of strong carbideformers (% Ta, % Zr, % Hf). All values are given in weight percentages.

Five examples show the properties that can be obtained:

Max. Hard./ Therm. diff. Tem. T° Hardness mm²/s at % C % Mo % W % V % Co% Mn Other HRc/° C. HRc 600° C. 0.32 3.36 1.52 0.45 2.66 <0.01 Hf, Zr,55/600 51.5 6.05 Nb 0.32 3.36 1.52 0.45 2.66 <0.01 Hf, Zr, 55/600 396.37 Nb 0.36 3.75 1.91 0.44 2.44 0.47 Hf, Zr, 57/600 53 6.03 Si = 0.20.34 4.04 1.23 0.73 2.16 0.6 Hf, Zr, 56/600 41 6.14 Nb 0.37 3.64 1.210.49 1.6 <0.01 Hf, Zr, 55/605 42 6.04 Ni = 2.7 0.51 3.75 1.51 <0.01 2.1<0.01 Hf, Zr, 51/600 44 6.42 Ni = 2.9 0.36 3.28 0.91 0.55 3.1 0.58 Hf,Zr, 56.5/610   38 6.83 0.61 3.6 1.19 0.56 2.6 0.54 Hf, Zr, 59/615 40.56.55 0.43 3.22 0.96 0.04 2.8 0.5 Hf, Zr, 56/600 47 6.26 0.32 3.25 0.960.43 2.45 0.41 Hf, Zr 56/610 48 6.34 0.33 3.48 0.86 <0.01 2.49 0.16 Hf,Zr 54/605 43 6.52 * In all cases heat treatment which maximizesdiffusivity at the indicated hardness has been applied, minimizing thepresence of elements in solution with the matrix, especially minimizingthe presence of % C and, to a lesser extent, % V in the matrix. In allcases this means high austenitization temperatures, from 3 to 5tempering cycles, with the latest in the range 640-690° C. Being % Hf:0.02-0.16, % Zr: 0.05-0.18 y % Nb: about 0.07, unless specificallyindicated.

Example 4

For hot cutting of sheet. In this case wear resistance must be maximizedwith a good hardenability and toughness (fracture toughness, in thiscase). Thermal conductivity is very important to maintain thetemperature at the cutting edge as low as possible. Weldability is lessimportant in this case, and small inserts are often used, socompositions with high content of alloying elements can be used. Forthis purpose, in the context of the present invention, the followingcompositional range can be used:

C_(eq): 0.5-0.9 Cr<0.1% (preferably Cr<0.05%)Si: <0.15% (preferably Si<0.1%)V: 0-2% for cases with Mo_(eq)>5 and V: 0-4% for cases with Mo_(eq)<5

Mo_(eq): 5-10

where

Mo_(eq)=% Mo+½% W

The rest of the elements should be kept as low as possible and, in anycase, always be below 0.15%, with the exception of strong carbideformers (% Ta, % Zr, % Hf). All values are given in weight percentages.

Three examples show the properties that can be obtained:

Therm. Diff. Hardness mm²/s at % C % Mo % W % V % Cr % Si % Mn Other HRc400° C. 0.59 6.7 4.6 <0.01 <0.01 <0.05 <0.01 — 55 12.39 0.69 7.89 3.950.7 <0.01 <0.05 <0.01 — 55 10.76 0.62 8.01 3.75 0.1 <0.01 <0.05 <0.01 Ni= 0.28 57 10.10 0.75 6.11 3.4 0.5 <0.01 <0.05 <0.01 Hf = 0.28 61 9.87 Zr= 0.14 0.87 6.92 4.4 0.7 <0.01 <0.05 <0.01 Hf = 0.23 64 9.03 Zr = 0.15 *In all cases heat treatment which maximizes diffusivity at the indicatedhardness has been applied, minimizing the presence of elements insolution with the matrix, especially minimizing the presence of % C and,to a lesser extent, % V in the matrix. In this case, also seeking thehighest possible presence of primary carbides. In all cases this meansvery high austenitization temperatures (1120° C. in the first two casesand 1240° C. in the last one), from 2 to 4 tempering cycles, with thelatest in the range 600-640° C.

1. A steel, in particular a hot work tool steel, with the followingcomposition, all percentages being indicated in weight percent: % C_(eq)= 0.3-0.9 % C = 0.3-0.9 % N = 0-0.6 % B = 0-0.6 % Cr <2.8 % Ni = 0-3.8 %Si = 0-1.4 % Mn = 0-3 % A1 = 0-2.5 % Mo = 0-10 % W = 0-10 % Ti = 0-2 %Ta = 0-3 % Zr = 0-3 % Hf = 0-3 % V = 0-4 % Nb = 0-1.5 % Cu = 0-2 % Co =0-6 % S = 0-1 % Se = 0-1 % Te = 0-1 % Bi = 0-1 % As = 0-1 % Sb = 0-1 %Ca = 0-1

the rest consisting of iron and unavoidable impurities, wherein% C_(eq)=% C+0.86*% N+1.2*% B,characterized in that% Mo+½·% W>3.0.
 2. A steel according to claim 1, wherein:when % C_(eq) is <0.35,then K>0.75,orwhen % C_(eq) is >=0.35,then K>0.84,orwhen % C_(eq) is >=0.35,then % Hf+% Zr+% Ta+% Nb>=0.01,being:K=% C_(eq)/(0.4+(% Mo_(eq(real))−4)*0.04173),and% Mo_(eq(real))=% Mo+0.52*% W.
 3. A steel according to claim 1, wherein:% Mo_(eq(real))>3.3%.
 4. A steel according to claim 1, wherein:% V+% Nb+% Hf+% Zr>0.1.
 5. A steel according to claim 1, wherein:% V+% Nb+% Hf+% Zr>1.2.
 6. A steel according to claim 1, wherein:% C_(eq) is >0.32 and % C>0.32.
 7. A steel according to claim 1,wherein:% C_(eq)>0.36.
 8. A steel according to claim 1, wherein:% C>0.4.
 9. A steel according to claim 1, wherein:% Mo+½·% W<10.0.
 10. A steel according to claim 1, wherein:% Mo+½·% W<4.5 with % Mo=0-4.5 and % W=0-9.
 11. A steel according toclaim 1, with the proviso that:when % C_(eq)<0.35,then % V<1.7.
 12. A steel according to claim 1,wherein:% V<1.8
 13. A steel according to claim 1, wherein:% Nb<0.09.
 14. A steel according to claim 1, wherein:% Ni<2.99.
 15. A steel according to claim 1, wherein:% Ni<1.0.
 16. A steel according to claim 1, wherein:when % Cr>2,then % Nb+% Ta+% Zr+% Hf>0.2.
 17. A steel according toclaim, wherein:% C_(eq)>0.32,% Mo_(eq)>3.2 and,% Cr<2.5,with the proviso that:when C_(eq)<=0.36 then:3.56<% Mo_(eq)/% C_(eq)<11.5 orwhen 0.36<C_(eq)<=0.38 then:3.56<% Mo_(eq)/% C_(eq)<14 orwhen 0.38<C_(eq) then:3.56<% Mo_(eq)/% C_(eq)<16.8,being% Mo_(eq)=% Mo+½·% W.
 18. A steel according to claim 1, wherein:% C_(eq)>=0.33 and K<0.81,beingK=% C_(eq)/(0.4+(% Mo_(eq(real))−4)*0.04173),and% Mo_(eq(real))=% Mo+0.52*% W.
 19. A steel according to claim 1 wherein,when subjected to a martensitic, bainitic or martensitic-bainitic quenchwith at least one tempering cycle at temperature above 590° C., ahardness above 47 HRc is obtainable with a low scattering structurecharacterized by a diffusivity of 9 mm²/s or more.
 20. A steel accordingto claim 1 wherein, when subjected to at least one tempering cycle attemperature 590° C., a hardness of 53 HRc or more is obtainable with alow scattering structure characterized by a thermal diffusivity above 9mm²/s.
 21. A steel according to claim 1 wherein:% C>0.32,% Co>1.3 and% V<2.8.
 22. A steel according to claim 21 wherein, when subjected to atleast one tempering cycle at temperature above 660° C., a hardness of 50HRc or more is obtainable with a low scattering structure characterizedby a diffusivity of 5.8 mm²/s or more at 600° C.
 23. A die, tool orpiece at least partially comprising a tool steel according to claim 1.24. A process to manufacture a hot work tool steel, characterised inthat a steel according to claim 1 to is subjected to a martensitic,bainitic or martensitic-bainitic quench with at least one temperingcycle at temperature above 590° C., so that a steel having a hardnessabove 47 HRc with a low scattering structure characterized by adiffusivity of 9 mm²/s or more is obtainable.
 25. A process tomanufacture a hot work tool steel according to claim 24, wherein a steelhaving a hardness above 53 HRc with a low scattering structurecharacterized by a diffusivity of 9 mm²/s or more is obtainable.
 26. Aprocess to manufacture a hot work tool steel according to claim 24,wherein the steel is subjected to at least one tempering cycle attemperature above 660° C., so that a steel having a hardness of 50 HRcor more with a low scattering structure characterized by a diffusivityof 5.8 mm²/s or more at 600° C. is obtainable.